EP4361304A1 - Stahlblech, element und verfahren zur herstellung des stahlblechs und des besagten elements - Google Patents

Stahlblech, element und verfahren zur herstellung des stahlblechs und des besagten elements Download PDF

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EP4361304A1
EP4361304A1 EP22864236.9A EP22864236A EP4361304A1 EP 4361304 A1 EP4361304 A1 EP 4361304A1 EP 22864236 A EP22864236 A EP 22864236A EP 4361304 A1 EP4361304 A1 EP 4361304A1
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steel sheet
steel
temperature
content
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English (en)
French (fr)
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EP4361304A4 (de
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Yoichiro Matsui
Tadachika CHIBA
Shinjiro Kaneko
Akinori Nakamura
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JFE Steel Corp
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JFE Steel Corp
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0273—Final recrystallisation annealing
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
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    • C22C38/008—Ferrous alloys, e.g. steel alloys containing tin
    • C—CHEMISTRY; METALLURGY
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
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    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/20—Ferrous alloys, e.g. steel alloys containing chromium with copper
    • C—CHEMISTRY; METALLURGY
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
    • C23C2/0224—Two or more thermal pretreatments
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
    • C23C2/06—Zinc or cadmium or alloys based thereon
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C30/00—Coating with metallic material characterised only by the composition of the metallic material, i.e. not characterised by the coating process
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/004—Dispersions; Precipitations
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y02—TECHNOLOGIES OR APPLICATIONS FOR MITIGATION OR ADAPTATION AGAINST CLIMATE CHANGE
    • Y02P—CLIMATE CHANGE MITIGATION TECHNOLOGIES IN THE PRODUCTION OR PROCESSING OF GOODS
    • Y02P10/00—Technologies related to metal processing
    • Y02P10/20—Recycling

Definitions

  • the present invention relates to a steel sheet for use in various applications including automobiles and home appliances, to a member using the steel sheet, and to methods for manufacturing the same.
  • 980-1180 MPa grade high strength steel sheets are increasingly applied to automobile frame parts and seat parts.
  • 980-1180 MPa grade high strength steel sheets tends to result in press cracking due to low ductility and low stretch flange formability.
  • these high strength steel sheets are required to be improved in formability compared to the conventional level.
  • weight reduction and increased rigidity of parts are being sought through utilization of laser welding.
  • Patent Literature 1 discloses that a high-ductility steel sheet having a TS (tensile strength) of 80 kgf/mm 2 or more and TS ⁇ El ⁇ 2500 kgf/mm 2 ⁇ % is obtained by annealing a steel that contains C: 0.10 to 0.45%, Si: 0.5 to 1.8%, and Mn: 0.5 to 3.0%, and aging the steel at 350 to 500°C for 1 to 30 minutes to form retained ⁇ .
  • TS tensile strength
  • Patent Literature 2 discloses that a steel sheet excellent in ductility (El) and stretch flange formability ( ⁇ ) is obtained by annealing a steel that contains C: 0.10 to 0.25%, Si: 1.0 to 2.0%, and Mn: 1.5 to 3.0%, cooling the steel to 450 to 300°C at 10°C/s or more, and holding the steel for 180 to 600 seconds, thereby controlling the space factor of retained austenite to 5% or more, the space factor of bainitic ferrite to 60% or more, and the space factor of polygonal ferrite to 20% or less.
  • Patent Literature 3 discloses that a steel sheet having a specific chemical composition is annealed, cooled to a range of temperatures of 150 to 350°C, and subsequently reheated to and held at 350 to 600°C.
  • the microstructure obtained as described above includes ferrite, tempered martensite, and retained austenite, and high ductility and high stretch flange formability can be imparted to the steel sheet.
  • This technique utilizes the so-called Q & P (quenching & partitioning) principle (quenching and partitioning of carbon from martensite to austenite) in which the steel in the cooling process is cooled once to a range of temperatures between the martensite start temperature (Ms temperature) and the martensite finish temperature (Mf temperature), and is subsequently reheated and held to stabilize retained ⁇ .
  • Q & P quenching & partitioning
  • this principle is applied to the development of high strength steels with high ductility and high stretch flange formability.
  • Patent Literature 4 discloses a technique that improves the Q & P process described above. Specifically, this technique achieves high ductility and high stretch flange formability by annealing a steel with a specific chemical composition at a temperature equal to or higher than Ae 3 temperature - 10°C to ensure that polygonal ferrite will be 5% or less, and cooling the steel to a relatively high finish cooling temperature of Ms - 10°C to Ms - 100°C to ensure that upper bainite will be formed when the steel is reheated to 350 to 450°C.
  • Patent Literature 5 discloses a technique that utilizes low-temperature bainite and high-temperature bainite to obtain a steel sheet with excellent ductility and low-temperature toughness. Specifically, a steel that contains C: 0.10 to 0.5% is annealed, cooled to 150 to 400°C at a cooling rate of 10°C/s or more, held at the above temperature for 10 to 200 seconds to form low-temperature bainite, then reheated to a range of temperatures of more than 400°C and 540°C or below, and held for 50 seconds or more to form high-temperature bainite. A steel sheet with excellent ductility and low-temperature toughness is thus obtained.
  • Patent Literature 6 discloses that a steel sheet having excellent ductility, flangeability, and weldability is obtained by annealing a steel sheet that contains C: 0.01 to 0.3%, Si: 0.005 to 2.5%, Mn: 0.01 to 3%, Mo: 0.01 to 0.3%, and Nb: 0.001 to 0.1% in a high-temperature region where the microstructure is almost a ⁇ single phase, cooling the steel sheet to a temperature range of 200 to 450°C, and holding the steel sheet at the temperature to ensure that the microstructure contains 50 to 97% bainite or bainitic ferrite as the main phase and 3 to 50% austenite as the second phase.
  • Patent Literature 1 The conventional TRIP steel described in Patent Literature 1 has excellent ductility, but its stretch flange formability is very low.
  • the microstructure is mainly bainitic ferrite and includes a small amount of ferrite. Because of this configuration, the steel sheet is excellent in stretch flange formability but is not necessarily high in ductility. Thus, further improvements in ductility are desired for application to parts that are difficult to form.
  • Patent Literature 3 realizes relatively high ductility and high stretch flange formability compared to the conventional TRIP steels and steels making use of bainitic ferrite.
  • fracture occurs in the formation of hard-to-form parts, such as center pillars, and further enhancements in ductility are desired.
  • a steel sheet according to this technique does not necessarily have a sufficient amount of uniform deformation that indicates the degree of resistance to fracture.
  • the amount of uniform deformation is a ductility indicator El that indicates the amount of elongation to the onset of necking and is written as U. El. A further increase in U. El is desired.
  • Patent Literature 4 cannot ensure sufficient ductility because the polygonal ferrite formation is lessened in order to reduce the amount of massive martensite. Furthermore, because the finish cooling temperature is set relatively high in order to enhance El, a large amount of non-transformed ⁇ remains at the termination of cooling and consequently massive martensite tends to remain.
  • Patent Literature 5 makes use of low-temperature bainite and high-temperature bainite in order to enhance ductility. Bainite that transforms at low temperatures has only a small contribution to ductility enhancement, and the use of high-temperature bainite tends to be accompanied by massive microstructures that remain. It is therefore difficult to impart high ductility and high stretch flange formability at the same time.
  • Patent Literature 6 offers improvements of materials having higher strength. While the technique is effective when the laser welding speed is fast and the heat input is small, significant HAZ softening occurs under normal-speed or low-speed welding conditions where the laser output is 4 to 6 kW and the welding speed is 3 to 5 mpm. Specifically, the joint is fractured at HAZ when subjected to a tensile test along the tensile axis perpendicular to the laser weld line. Furthermore, the ductility is not necessarily high, and further enhancements in ductility are desired.
  • the present invention has been made to solve the problems discussed above and provides a steel sheet that has 980 MPa or higher tensile strength, has high ductility and excellent stretch flange formability, and further has excellent laser weldability; a related member; and methods for manufacturing the same.
  • steel sheet includes galvanized steel sheet resulting from surface galvanizing treatment.
  • 980 MPa or higher tensile strength means that the tensile strength in accordance with JIS Z2241 is 980 MPa or more.
  • the term high ductility means that the total elongation T-El in accordance with JIS Z2241 is 16.0% or more when TS is less than 1180 MPa, 14.0% or more when TS is 1180 MPa or more and less than 1320 MPa, and 13.0% or more when TS is 1320 MPa or more.
  • d 0 initial hole diameter (mm)
  • d hole diameter (mm) at the occurrence of cracking.
  • a 100 mm ⁇ 100 mm square sample is punched with a punching tool having a punch diameter of 10 mm and a die diameter of 10.3 mm (13% clearance), and a conical punch having an apex angle of 60 degrees is inserted into the hole in such a manner that the burr produced at the time of punching will be directed to the outside.
  • the hole is expanded until the sheet is cracked through the thickness.
  • the term excellent laser weldability means that a specimen obtained by laser welding described below is fractured in the base material fracture mode in a fracture mode determination test described below and satisfies HAZ strength ⁇ base material TS + 50 MPa in a notched tensile test described below.
  • the present invention has been made based on the above findings. Specifically, the present invention provides the following.
  • a steel sheet having high ductility, excellent stretch flange formability, and excellent laser weldability, and a related member can be obtained. Furthermore, strength can be increased according to the present invention.
  • the steel sheet of the present invention may be applied to an automotive part to realize weight reduction of the automotive part. Enhanced fuel efficiency is thus expected.
  • a steel sheet of the present invention has a chemical composition including, in mass%, C: 0.06 to 0.25%, Si: 0.4 to 2.5%, Mn: 1.5 to 3.5%, P: 0.02% or less, S: 0.01% or less, sol. Al: less than 1.0%, and N: less than 0.015%, the balance being Fe and incidental impurities.
  • the steel sheet includes a steel microstructure including, in area fraction, polygonal ferrite: 10% or less (including 0%), tempered martensite: 40% or more, fresh martensite: 20% or less (including 0%), bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 : 3 to 40%, and, in volume fraction, retained austenite: 5 to 20%.
  • the steel sheet has a proportion Sc ⁇ 0.5 /Sc ⁇ 0.3 ⁇ 100 of 20% or more wherein S C ⁇ 0.5 is the area of a region having a C concentration of 0.50% or more and S C ⁇ 0.3 is the area of a region having a C concentration of 0.30% or more.
  • the steel sheet of the present invention will be described below in the order of its chemical composition and its steel microstructure.
  • the steel sheet of the present invention includes the components described below.
  • the unit “%” for the contents of components means “mass%”.
  • Carbon is added to ensure predetermined strength by ensuring an area fraction of tempered martensite, to enhance ductility by ensuring a volume fraction of retained ⁇ , and to stabilize retained ⁇ by being concentrated in retained ⁇ and thereby to enhance ductility. Furthermore, the addition of carbon increases the strength of a fused portion of a welded joint and a portion quenched from ⁇ region, and thereby can eliminate or reduce deformation occurring at HAZ and enhance the HAZ softening resistance. When the C content is less than 0.06%, these effects cannot be ensured sufficiently. Thus, the lower limit is limited to 0.06%.
  • the C content is preferably 0.09% or more, and more preferably 0.11% or more.
  • the C content exceeds 0.25%, upper bainite transformation during intermediate holding in the course of cooling is retarded, and it becomes difficult to form a predetermined amount of retained ⁇ that is adjacent to upper bainite. As a result, ductility is lowered. Furthermore, the amount of massive martensite or massive retained ⁇ is increased to deteriorate stretch flange formability. Furthermore, laser welding characteristics of the steel sheet, such as HAZ softening resistance, spot weldability, bendability, and flangeability, are significantly deteriorated. Thus, the upper limit of the C content is limited to 0.25%. From the points of view of ductility and HAZ softening resistance, the C content is preferably 0.22% or less. In order to further improve ductility and HAZ softening resistance, the C content is more preferably 0.20% or less.
  • the Si content is limited to 0.4% or more.
  • the Si content is preferably 0.6% or more. More preferably, the Si content is 0.8% or more.
  • the Si content exceeds 2.5%, the rolling load at the time of hot rolling is extremely increased to make sheet production difficult.
  • chemical convertibility and weld toughness are deteriorated.
  • the Si content is limited to 2.5% or less.
  • the Si content is preferably less than 2.0%.
  • the Si content is preferably 1.8% or less, and more preferably 1.5% or less.
  • Manganese is an important element from the points of view of ensuring strength by ensuring a predetermined area fraction of tempered martensite and/or bainite; improving ductility by lowering the Ms temperature of retained ⁇ and thereby stabilizing retained ⁇ ; enhancing ductility by suppressing the formation of carbides in bainite similarly to silicon; and enhancing ductility by increasing the volume fraction of retained ⁇ .
  • the Mn content is limited to 1.5% or more.
  • the Mn content is preferably 2.5% or more.
  • the Mn content is preferably 2.6% or more, and more preferably 2.7% or more.
  • the Mn content exceeds 3.5%, bainite transformation is significantly retarded to lower ductility and HAZ softening resistance.
  • the Mn content exceeds 3.5%, moreover, it becomes difficult to suppress the formation of massive coarse ⁇ and massive coarse martensite, resulting in a decrease in stretch flange formability.
  • the Mn content is limited to 3.5% or less.
  • the Mn content is preferably 3.2% or less. More preferably, the Mn content is 3.1% or less.
  • Phosphorus is an element that strengthens steel, but much phosphorus deteriorates spot weldability.
  • the P content is limited to 0.02% or less.
  • the P content is preferably 0.01% or less.
  • the P content may be nil. From the point of view of manufacturing cost, the P content is preferably 0.001% or more.
  • Sulfur is an element that is effective in improving scale exfoliation in hot rolling and effective in suppressing nitridation during annealing, but sulfur lowers spot weldability, bendability, and flangeability.
  • the S content is limited to 0.01% or less.
  • the contents of C, Si, and Mn are high and spot weldability tends to be lowered.
  • the S content is preferably 0.0020% or less, and more preferably less than 0.0010%.
  • the S content may be nil. From the point of view of manufacturing cost, the S content is preferably 0.0001% or more. More preferably, the S content is 0.0005% or more.
  • Aluminum is added for the purposes of deoxidization and stabilizing retained ⁇ as a substitute for silicon.
  • the lower limit of the sol. Al content is not particularly limited.
  • the sol. Al content is preferably 0.005% or more.
  • the sol. Al content is more preferably 0.01% or more.
  • 1.0% or more sol. Al significantly lowers the strength of the base material and also affects adversely chemical convertibility.
  • the sol. Al content is limited to less than 1.0%.
  • the sol. Al content is more preferably less than 0.50%, and still more preferably 0.20% or less.
  • Nitrogen is an element that forms nitrides, such as BN, AlN, and TiN, in steel. This element lowers the hot ductility of steel and lowers the surface quality. Furthermore, in B-containing steel, nitrogen has a harmful effect in eliminating the effect of boron through the formation of BN. The surface quality is significantly deteriorated when the N content is 0.015% or more. Thus, the N content is limited to less than 0.015%. The N content is preferably 0.010% or less. The N content may be nil. From the point of view of manufacturing cost, the N content is preferably 0.0001% or more. More preferably, the N content is 0.001% or more.
  • the balance after the above components is Fe and incidental impurities.
  • the steel sheet of the present invention preferably has a chemical composition that contains the basic components described above, with the balance consisting of iron (Fe) and incidental impurities.
  • the chemical composition of the steel sheet of the present invention may appropriately include one, or two or more optional elements selected from the following (A), (B), and (C):
  • Titanium fixes nitrogen in steel as TiN to produce an effect of enhancing hot ductility and an effect of allowing boron to produce its effect of enhancing hardenability. Furthermore, titanium has an effect of reducing the size of the microstructure through TiC precipitation. In order to obtain these effects, the Ti content is preferably 0.002% or more. In order to fix nitrogen sufficiently, the Ti content is more preferably 0.008% or more. The Ti content is still more preferably 0.010% or more. On the other hand, more than 0.1% titanium may cause an increase in rolling load and a decrease in ductility by an increased amount of precipitation strengthening. Thus, when titanium is added, the Ti content is limited to 0.1% or less. Preferably, the Ti content is 0.05% or less. In order to ensure high ductility, the Ti content is more preferably 0.03% or less.
  • Boron is an element that enhances the hardenability of steel and facilitates the formation of a predetermined area fraction of tempered martensite and/or bainite. Furthermore, boron enhances the hardenability in the vicinity of a weld and allows a hard microstructure to be formed in the vicinity of the weld, thereby enhancing HAZ softening resistance. Furthermore, residual solute boron enhances delayed fracture resistance.
  • the B content is preferably 0.0002% or more.
  • the B content is more preferably 0.0005% or more. Still more preferably, the B content is 0.0010% or more.
  • the B content exceeds 0.01%, the effects are saturated, and further hot ductility is significantly lowered to invite surface defects.
  • the B content is limited to 0.01% or less.
  • the B content is 0.0050% or less. More preferably, the B content is 0.0030% or less.
  • Copper enhances the corrosion resistance in automobile use environments. Furthermore, corrosion products of copper cover the surface of the steel sheet and can suppress penetration of hydrogen into the steel sheet. Copper is an element that is mixed when scraps are used as raw materials. By accepting copper contamination, recycled materials can be used as raw materials and thereby manufacturing costs can be reduced. From these points of view, the Cu content is preferably 0.005% or more, and, further from the point of view of enhancing delayed fracture resistance, the Cu content is more preferably 0.05% or more. Still more preferably, the Cu content is 0.10% or more. On the other hand, too much copper invites surface defects. Thus, when copper is added, the Cu content is limited to 1% or less. The Cu content is preferably 0.4% or less, and more preferably 0.2% or less.
  • nickel can enhance corrosion resistance. Furthermore, nickel can also eliminate or reduce the occurrence of surface defects that tend to occur when the steel contains copper. To benefit from these effects, it is preferable to add 0.01% or more nickel.
  • the Ni content is more preferably 0.04% or more, and still more preferably 0.06% or more.
  • adding too much nickel can instead cause surface defects because scales are formed nonuniformly in a heating furnace, and also increases the cost.
  • the Ni content is limited to 1% or less.
  • the Ni content is preferably 0.4% or less, and more preferably 0.2% or less.
  • Chromium may be added to produce an effect of enhancing the hardenability of steel and an effect of suppressing the formation of carbides in martensite and upper/lower bainite. Furthermore, chromium enhances the hardenability in the vicinity of a weld and allows a hard phase to be formed in the vicinity of the weld, thereby enhancing HAZ softening resistance.
  • the Cr content is preferably 0.01% or more.
  • the Cr content is more preferably 0.03% or more, and still more preferably 0.06% or more.
  • too much chromium deteriorates pitting corrosion resistance.
  • the Cr content is limited to 1.0% or less.
  • the Cr content is preferably 0.8% or less, and more preferably 0.4% or less.
  • Molybdenum may be added to produce an effect of enhancing the hardenability of steel and an effect of suppressing the formation of carbides in martensite and upper/lower bainite. Furthermore, molybdenum enhances the hardenability in the vicinity of a weld and allows a hard phase to be formed in the vicinity of the weld, thereby enhancing HAZ softening resistance.
  • the Mo content is preferably 0.01% or more.
  • the Mo content is more preferably 0.03% or more, and still more preferably 0.06% or more.
  • molybdenum significantly deteriorates the chemical convertibility of the cold rolled steel sheet. Thus, when molybdenum is added, the Mo content is limited to 0.5% or less. From the point of view of enhancing chemical convertibility, the Mo content is preferably 0.15% or less.
  • Vanadium may be added to produce an effect of enhancing the hardenability of steel, an effect of suppressing the formation of carbides in martensite and upper/lower bainite, an effect of reducing the size of the microstructure, and an effect of improving delayed fracture resistance through the precipitation of carbide. Furthermore, vanadium enhances the hardenability in the vicinity of a weld and allows a hard phase to be formed in the vicinity of the weld, thereby enhancing HAZ softening resistance. In order to obtain these effects, the V content is preferably 0.003% or more. The V content is more preferably 0.005% or more, and still more preferably 0.010% or more. On the other hand, much vanadium significantly deteriorates castability.
  • the V content when vanadium is added, the V content is limited to 0.5% or less.
  • the V content is preferably 0.3% or less, and more preferably 0.1% or less.
  • the V content is still more preferably 0.05% or less, and further preferably 0.03% or less.
  • Niobium may be added to produce an effect of reducing the size of the steel microstructure and thereby increasing the strength, and, through grain size reduction, an effect of promoting bainite transformation, an effect of improving bendability, and an effect of enhancing delayed fracture resistance. Furthermore, niobium enhances the hardenability in the vicinity of a weld and allows a hard phase to be formed in the vicinity of the weld, thereby enhancing HAZ softening resistance.
  • the Nb content is preferably 0.002% or more.
  • the Nb content is more preferably 0.004% or more, and still more preferably 0.010% or more.
  • adding much niobium results in excessive precipitation strengthening and low ductility. Furthermore, the rolling load is increased and castability is deteriorated. Thus, when niobium is added, the Nb content is limited to 0.1% or less.
  • the Nb content is preferably 0.05% or less, and more preferably 0.03% or less.
  • Zirconium may be added to produce an effect of enhancing the hardenability of steel, an effect of suppressing the formation of carbides in bainite, an effect of reducing the size of the microstructure, and an effect of improving delayed fracture resistance through the precipitation of carbide.
  • the Zr content is preferably 0.005% or more.
  • the Zr content is more preferably 0.008% or more, and still more preferably 0.010% or more.
  • the steel contains much zirconium, increased amounts of coarse precipitates, such as ZrN and ZrS, remain undissolved at the time of slab heating before hot rolling to cause deterioration in delayed fracture resistance.
  • the Zr content is limited to 0.2% or less.
  • the Zr content is preferably 0.15% or less, and more preferably 0.08% or less.
  • the Zr content is still more preferably 0.03% or less, and further preferably 0.02% or less.
  • Tungsten may be added to produce an effect of enhancing the hardenability of steel, an effect of suppressing the formation of carbides in bainite, an effect of reducing the size of the microstructure, and an effect of improving delayed fracture resistance through the precipitation of carbide.
  • the W content is preferably 0.005% or more.
  • the W content is more preferably 0.008% or more, and still more preferably 0.010% or more.
  • the steel contains much tungsten, increased amounts of coarse precipitates, such as WN and WS, remain undissolved at the time of slab heating before hot rolling to cause deterioration in delayed fracture resistance.
  • the W content is limited to 0.2% or less.
  • the W content is preferably 0.15% or less, and more preferably 0.08% or less.
  • the W content is still more preferably 0.03% or less, and further preferably 0.02% or less.
  • the Ca content is preferably 0.0002% or more.
  • the Ca content is more preferably 0.0005% or more, and still more preferably 0.0010% or more.
  • much calcium deteriorates surface quality and bendability.
  • the Ca content is limited to 0.0040% or less.
  • the Ca content is preferably 0.0035% or less, and more preferably 0.0020% or less.
  • the Ce content is preferably 0.0002% or more.
  • the Ce content is more preferably 0.0004% or more, and still more preferably 0.0006% or more.
  • much cerium deteriorates surface quality and bendability.
  • the Ce content is limited to 0.0040% or less.
  • the Ce content is preferably 0.0035% or less, and more preferably 0.0020% or less.
  • the La content is preferably 0.0002% or more.
  • the La content is more preferably 0.0004% or more, and still more preferably 0.0006% or more.
  • much lanthanum deteriorates surface quality and bendability.
  • the La content is limited to 0.0040% or less.
  • the La content is preferably 0.0035% or less, and more preferably 0.0020% or less.
  • the Mg content is preferably 0.0002% or more.
  • the Mg content is more preferably 0.0004% or more, and still more preferably 0.0006% or more.
  • much magnesium deteriorates surface quality and bendability.
  • the Mg content is limited to 0.0030% or less.
  • the Mg content is preferably 0.0025% or less, and more preferably 0.0010% or less.
  • the Sb content is preferably 0.002% or more.
  • the Sb content is more preferably 0.004% or more, and still more preferably 0.006% or more.
  • the Sb content exceeds 0.1%, castability is deteriorated and segregation occurs at prior ⁇ grain boundaries to deteriorate the delayed fracture resistance of sheared end faces.
  • the Sb content is limited to 0.1% or less.
  • the Sb content is preferably 0.04% or less, and more preferably 0.03% or less.
  • Tin suppresses oxidation and nitridation of a superficial portion of the steel sheet and thereby eliminates or reduces the loss of the C and B contents in the superficial portion. Furthermore, the elimination or reduction of the loss of the C and B contents leads to suppressed formation of ferrite in the superficial portion of the steel sheet, thus increasing strength and improving delayed fracture resistance.
  • the Sn content is preferably 0.002% or more.
  • the Sn content is preferably 0.004% or more, and more preferably 0.006% or more.
  • the Sn content exceeds 0.1%, castability is deteriorated.
  • tin is segregated at prior ⁇ grain boundaries to deteriorate the delayed fracture resistance of sheared end faces. Thus, when tin is added, the Sn content is limited to 0.1% or less.
  • the Sn content is preferably 0.04% or less, and more preferably 0.03% or less.
  • Polygonal ferrite 10% or less (including 0%)
  • Polygonal ferrite which is formed during annealing or a cooling process, contributes to enhancement in ductility but decreases stretch flange formability by giving rise to a difference in hardness from surrounding hard phases, such as martensite. Polygonal ferrite does not impair the advantageous effects of the present invention as long as the area fraction thereof is 10% or less and therefore, polygonal ferrite may be contained up to 10% or less in area fraction. Thus, in the present invention, the area fraction of polygonal ferrite is limited to 10% or less.
  • the polygonal ferrite is preferably 5% or less, and more preferably 2% or less.
  • the polygonal ferrite may be 0%.
  • Tempered martensite 40% or more
  • the area fraction of tempered martensite is limited to 40% or more.
  • the tempered martensite is preferably 50% or more.
  • the tempered martensite exceeds 80%, the strength is excessively increased and the ductility is lowered.
  • the tempered martensite is preferably 80% or less.
  • the tempered martensite is more preferably 75% or less.
  • Fresh martensite 20% or less (including 0%)
  • the final tempering step (the residence step at an average cooling rate CR4 described later) to produce a large amount of bainite transformation conventionally results in a large amount of massive martensite or massive retained ⁇ that remains.
  • the conventional approach to preventing this is to reduce the amount of manganese and thereby to promote bainite transformation.
  • decreasing the Mn content lowers ductility because of the loss of the effects of stabilizing retained ⁇ and increasing the volume fraction of retained ⁇ .
  • the present invention performs an appropriate cooling treatment on the steel sheet containing a large amount of manganese, and thereby can make use of bainite transformation and reduce the occurrence of massive microstructures at the same time.
  • Excellent stretch flange formability and HAZ softening resistance can be ensured by reducing the area fraction of massive fresh martensite microstructures to 20% or less.
  • the area fraction of fresh martensite in the present invention is limited to 20% or less.
  • the fresh martensite is preferably 10% or less.
  • the fresh martensite is more preferably 5% or less.
  • the fresh martensite may be 0%.
  • Bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 : 3 to 40%
  • bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 When bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 are present with an area fraction of 3% or more, carbon is efficiently concentrated into surrounding retained ⁇ . Furthermore, such bainitic ferrite is insusceptible to heat at the time of laser welding and contributes to enhancement in HAZ softening resistance.
  • the area fraction of bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 is limited to 3% or more.
  • the bainitic ferrite is preferably 5% or more, and more preferably 7% or more.
  • the area fraction of bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 In order to reduce the decrease in strength, the area fraction of bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 is limited to 40% or less.
  • the area fraction is preferably 30% or less, and more preferably 25% or less.
  • the steel may contain bainitic ferrite having more
  • Microstructure including one, or two or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite: 90% or more (including 100%)
  • the remaining microstructure after the polygonal ferrite have a total area fraction of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite of 90% or more.
  • the remaining microstructure may be a microstructure including one, or two or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite, or may be a microstructure consisting of one, or two or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite.
  • the upper bainite and the lower bainite include bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 .
  • the upper bainite and the lower bainite may include bainitic ferrite having more than 20 internal carbides per 10 ⁇ m 2 .
  • the volume fraction of retained austenite (retained ⁇ ) is limited to 5% or more of the entire steel microstructure.
  • the retained austenite is preferably 7% or more, and more preferably 9% or more.
  • This amount of retained ⁇ includes retained ⁇ formed adjacent to bainite.
  • An excessively large amount of retained ⁇ invites decreases in strength, stretch flange formability, and delayed fracture resistance.
  • the volume fraction of retained ⁇ is limited to 20% or less.
  • the retained austenite is preferably 15% or less.
  • the "volume fraction" may be regarded as the "area fraction".
  • the proportion S C ⁇ 0.5 /S C ⁇ 0.3 ⁇ 100 is limited to 20% or more.
  • S C ⁇ 0.5 is the area of a region having a C concentration of 0.50% or more
  • S C ⁇ 0.3 is the area of a region having a C concentration of 0.30% or more.
  • the proportion is preferably 25% or more, and more preferably 30% or more.
  • bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 carbon is efficiently partitioned to adjacent non-transformed austenite, and consequently retained ⁇ in the final microstructure attains a high carbon concentration and can contribute to enhancement in ductility.
  • the number density of retained austenite present adjacent to bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 be 50 or more per 10000 ⁇ m 2 . More preferably, the number density is 70 or more, and still more preferably 100 or more per 10000 ⁇ m 2 .
  • the number density of retained austenite present adjacent to bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 is preferably 400 or less, and more preferably 300 or less per 10000 ⁇ m 2 .
  • Fig. 2 illustrates an example SEM image of the steel microstructure of the steel sheet.
  • the polygonal ferrite discussed here is relatively equiaxed ferrite containing almost no internal carbides. This region looks blackest in SEM.
  • Bainitic ferrite is a ferrite microstructure that contains internal carbide or retained ⁇ which looks white in SEM.
  • the area fractions are calculated while assuming that the region is polygonal ferrite when the aspect ratio is ⁇ 2.0 and the region is bainitic ferrite when the aspect ratio is > 2.0.
  • Fig. 3 is a set of views illustrating a manner for measuring the steel microstructure of the steel sheet of the present invention. As illustrated in Fig.
  • the aspect ratio is obtained from a/b in which a is the major axis length where the particle length is longest, and b is the minor axis length that cuts across the particle over the largest length perpendicular to the major axis length.
  • a is the major axis length where the particle length is longest
  • b is the minor axis length that cuts across the particle over the largest length perpendicular to the major axis length.
  • the number of internal carbides in bainitic ferrite per 10 ⁇ m 2 can be determined by measuring the area fraction of bainitic ferrite and counting the number of internal carbides in a ⁇ 5000 SEM image, dividing the carbide count by the area fraction of bainitic ferrite, and converting the quotient into the value per 10 ⁇ m 2 .
  • Tempered martensite is a region that contains a lath-like submicrostructure and carbide precipitates according to SEM.
  • Fresh martensite is a massive region that looks white and does not contain any visible submicrostructures according to SEM.
  • the microstructure including one, or two or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite corresponds to the remaining microstructure after the polygonal ferrite, and the total area fraction of this microstructure is the area fraction of the regions other than the polygonal ferrite.
  • the area fraction of carbides is very small and is thus included in the above area fraction of the remaining microstructure.
  • the volume fraction of retained austenite (retained ⁇ ) was determined by chemically polishing the steel sheet surface to a location at 1/4 thickness and analyzing the sheet surface by X-ray diffractometry. Co-K ⁇ radiation source is used as the incident X-ray, and the volume fraction of retained austenite is calculated from the intensity ratio of (200), (211), and (220) planes of ferrite and (200), (220), and (311) planes of austenite. Because retained ⁇ is randomly distributed, the volume fraction of retained ⁇ obtained by X-ray diffractometry is equal to the area fraction of retained ⁇ in the steel microstructure.
  • the number density of retained austenite present adjacent to bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 is determined as follows.
  • the sample used for bainitic ferrite observation is mirror-polished.
  • An electron backscattering diffraction pattern (EBSD) of the same field of view as obtained in SEM is subjected to mapping measurement using EBSD analysis program OIM Data Collection ver. 7.
  • the obtained data is analyzed using TSL OIM Analysis ver. 7 (manufactured by EDAX/TSL) to give phase map data.
  • the number density of fcc structures present adjacent to bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 is measured.
  • adjacent means that the fcc structures are in contact with bcc structures corresponding to bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 in the phase map.
  • the fcc structure may be included in the bcc structure.
  • the area S C ⁇ 0.5 of a region having a C concentration of 0.50% (mass%) or more and the area S C ⁇ 0.3 of a region having a C concentration of 0.30% (mass%) or more are measured by mapping analysis of the C concentration distribution with respect to positions at 1/4 thickness of a through-thickness cross section perpendicular to the steel sheet surface and parallel to the rolling direction, using field emission electron probe microanalyzer (FE-EPMA) JXA-8500F manufactured by JEOL Ltd., at an acceleration voltage of 6 kV and an illumination current of 7 ⁇ 10 -8 A with the minimum beam diameter.
  • FE-EPMA field emission electron probe microanalyzer
  • the background is subtracted so that the average value of carbon obtained by the analysis will be equal to the amount of carbon in the base material. Specifically, when the average of the measured amounts of carbon is greater than the amount of carbon in the base material, the excess is understood as contamination, and the excess is subtracted from each of the values analyzed at the respective positions. The values thus obtained are taken as the true amounts of carbon at the respective positions.
  • the steel sheet of the present invention preferably has a tensile strength of 980 MPa or more. More preferably, the tensile strength is 1180 MPa or more.
  • the upper limit of the tensile strength is preferably 1450 MPa or less from the point of view of compatibility with other characteristics, and is more preferably 1400 MPa or less.
  • the total elongation T-El is 16.0% or more when TS is less than 1180 MPa, 14.0% or more when TS is 1180 MPa or more and less than 1320 MPa, and 13.0% or more when TS is 1320 MPa or more.
  • the hole expansion ratio ⁇ is 30% or more.
  • the upper limit of ⁇ is preferably 90% or less, and more preferably 80% or less at any level of strength.
  • the steel sheet of the present invention is preferably such that a specimen obtained by laser welding is fractured in the base material fracture mode in a fracture mode determination test and satisfies HAZ strength ⁇ base material TS + 50 MPa in a notched tensile test.
  • the steel sheet of the present invention described above may be a steel sheet having a galvanized layer on a surface.
  • the galvanized layer may be a hot-dip galvanized layer or an electrogalvanized layer.
  • a steel slab having the chemical composition described hereinabove is hot rolled and cold rolled.
  • the cold rolled steel sheet obtained is annealed.
  • the annealing includes the following steps in the order named: a step of holding the steel sheet at an annealing temperature of 810 to 900°C; a step of cooling the steel sheet in a range of temperatures from 810°C to 500°C at an average cooling rate (CR1) of 5 to 100°C/s; a step of causing the steel sheet to reside in a range of temperatures from 500°C to a residence finish temperature (T1) that is equal to or higher than a martensite start temperature Ms (°C) and is equal to or higher than 320°C for 10 seconds or more and 60 seconds or less while cooling the steel sheet at an average cooling rate (CR2) of 10°C/s or less; a step of cooling the steel sheet in a range of temperatures from the residence finish temperature (T1) to a finish cooling temperature (T2) of 200°C
  • the temperatures specified in the steps in the present invention indicate the surface temperatures of the slab (steel slab) or the steel sheet.
  • Fig. 4 is a diagram illustrating the method for manufacturing the steel sheet according to the present invention, in particular, indicating changes in surface temperature of the slab (steel slab) or the steel sheet with time. The details of the steps, including the changes in temperature with time, will be described below.
  • the steel slab may be hot rolled in such a manner that the slab is heated and then rolled, that the slab from continuous casting is subjected to hot direct rolling without heating, or that the slab from continuous casting is quickly heat treated and then rolled.
  • the hot rolling may be performed in accordance with a conventional procedure.
  • the slab heating temperature may be 1100 to 1300°C; the soaking time may be 20 to 300 minutes; the finish rolling temperature may be Ar 3 transformation temperature to Ar 3 transformation temperature + 200°C; and the coiling temperature may be 400 to 720°C.
  • the coiling temperature is preferably 430 to 530°C.
  • the rolling reduction ratio (the cumulative rolling reduction ratio) may be 30 to 85%. In order to ensure high strength stably and to reduce anisotropy, the rolling reduction ratio is preferably 35 to 85%.
  • a softening annealing treatment may be performed on CAL (a continuous annealing line) or in BAF (a box annealing furnace) at 450 to 730°C.
  • the steel sheet is annealed under the conditions specified below.
  • the annealing facility is not particularly limited, but a continuous annealing line (CAL) or a continuous galvanizing line (CGL) is preferable from the points of view of productivity and ensuring the desired heating rate and cooling rate.
  • CAL continuous annealing line
  • CGL continuous galvanizing line
  • the annealing temperature is limited to 810 to 900°C.
  • the annealing temperature is preferably adjusted so that the annealing will take place in the ⁇ single-phase region.
  • the annealing temperature is preferably 815°C or above.
  • the annealing temperature exceeds 900°C, the ⁇ grain size is excessively increased to extend the distance of diffusion of carbon atoms required to obtain retained ⁇ having the desired carbon concentration, resulting in a decrease in ductility.
  • the annealing temperature is limited to 900°C or below.
  • the annealing temperature is 880°C or below.
  • Cooling in a range of temperatures from 810°C to 500°C at an average cooling rate (CR1) of 5 to 100°C/s
  • the steel sheet is cooled in a range of temperatures from 810°C to 500°C at an average cooling rate (CR1) of 5 to 100°C/s.
  • the average cooling rate (CR1) is preferably 8°C/s or more.
  • the average cooling rate (CR1) is preferably 50°C/s or less, and more preferably less than 30°C/s.
  • the average cooling rate (CR1) is "(810°C (cooling start temperature) - 500°C (finish cooling temperature))/(cooling time (seconds) from cooling start temperature 810°C to finish cooling temperature 500°C)".
  • the steel sheet is caused to reside (is gradually cooled) in a range of temperatures from 500°C to a residence finish temperature (T1) that is equal to or higher than a martensite start temperature Ms (°C) and is equal to or higher than 320°C for 10 seconds or more and 60 seconds or less while being cooled at an average cooling rate (CR2) of 10°C/s or less.
  • T1 a residence finish temperature
  • Ms martensite start temperature
  • CR2 average cooling rate
  • the temperature range exceeds 500°C
  • the driving force for bainite transformation decreases and the amount of bainite transformation is reduced.
  • the temperature range is limited to be 500°C or below and to be equal to or higher than the Ms and equal to or higher than 320°C.
  • This temperature range is preferably 380°C or above, and more preferably 420°C or above.
  • the temperature range is preferably 480°C or below, and more preferably 460°C or below.
  • the average cooling rate (CR2) exceeds 10°C/s
  • the amount of bainite transformation is reduced.
  • the average cooling rate (CR2) is limited to 10°C/s or less.
  • the residence time is less than 10 seconds, the desired amount of bainite cannot be obtained.
  • the residence time exceeds 60 seconds, the enrichment of carbon from bainite to massive non-transformed ⁇ proceeds to result in an increase in the residual amount of the massive microstructure.
  • the residence time is limited to 10 seconds or more and 60 seconds or less.
  • the residence time is preferably 20 seconds or more.
  • the residence time is preferably 50 seconds or less.
  • the martensite start temperature Ms can be determined using a Formaster tester by holding a cylindrical test piece (3 mm in diameter ⁇ 10 mm in height) at a predetermined annealing temperature and quenching the test piece with helium gas while measuring the volume change.
  • the average cooling rate (CR2) is "(500°C (residence start temperature) - residence finish temperature (T1))/(residence time (seconds) from 500°C to residence finish temperature (T1))".
  • the steel sheet After the above residence, the steel sheet needs to be cooled rapidly to avoid excessive progress of carbon enrichment into ⁇ .
  • the average cooling rate (CR3) in the range of temperatures from the residence finish temperature T1 of 320°C or above to the finish cooling temperature T2 of 200°C or above and 300°C or below is less than 3°C/s, carbon is concentrated into massive non-transformed ⁇ and an increased amount of fresh martensite is formed during the final cooling to cause a decrease in stretch flange formability.
  • the average cooling rate (CR3) in the range of temperatures from the residence finish temperature T1 to the finish cooling temperature T2 of 200°C or above and 300°C or below is limited to 3°C/s or more.
  • the average cooling rate (CR3) is more preferably 5°C/s or more, and still more preferably 8°C/s or more.
  • the average cooling rate (CR3) in the above temperature range is limited to 100°C/s or less.
  • the average cooling rate (CR3) is preferably 50°C/s or less.
  • the finish cooling temperature T2 is limited to 200°C or above.
  • the finish cooling temperature T2 is preferably 220°C or above, and more preferably 240°C or above.
  • finish cooling temperature T2 When the finish cooling temperature T2 is above 300°C, a large amount of massive non-transformed ⁇ remains and an increased amount of fresh martensite is formed during the final cooling to cause a decrease in stretch flange formability. Thus, the finish cooling temperature T2 is limited to 300°C or below. The finish cooling temperature T2 is preferably 280°C or below.
  • the average cooling rate (CR3) is "(residence finish temperature (T1)) - (finish cooling temperature (T2))/(cooling time (seconds) from residence finish temperature (T1) to finish cooling temperature (T2))".
  • the steel sheet is heated in a range of temperatures from the finish cooling temperature (T2) to 380°C in a short time. In this manner, the precipitation of carbides is suppressed and high ductility can be ensured.
  • upper bainite is formed from martensite or bainite as nucleus that has been formed during cooling.
  • the average heating rate during heating to 380°C is low, the above effects cannot be obtained. As a result, the amount of retained ⁇ is reduced and ductility is lowered.
  • the average heating rate in the range of temperatures from the finish cooling temperature (T2) to 380°C is limited to 2°C/s or more.
  • the average heating rate is preferably 5°C/s or more, and more preferably 10°C/s or more.
  • the upper limit of the average heating rate is not particularly limited but is preferably 50°C/s or less, and more preferably 30°C/s or less.
  • the average heating rate is "380°C (heating stop temperature) - (finish cooling temperature (T2))/(heating time (seconds) from finish cooling temperature T2 to 380°C (heating stop temperature))".
  • the steel sheet In order to partition carbon to retained ⁇ and stabilize the retained ⁇ and in order to divide massive regions distributed as non-transformed ⁇ by bainite transformation and thereby to enhance stretch flange formability, the steel sheet is caused to reside (is gradually cooled) in a range of temperatures of 340°C or above and 590°C or below for 20 seconds or more and 3000 seconds or less. Furthermore, in order to enhance stretch flange formability by eliminating or reducing the formation of massive microstructures due to excessive partitioning of carbon to retained ⁇ and also by causing self-tempering of fresh martensite, the steel sheet is cooled slowly in this temperature range at an average cooling rate (CR4) of 0.01 to 5°C/s.
  • CR4 average cooling rate
  • the average cooling rate (CR4) is less than 0.01°C/s, carbon is excessively partitioned to retained ⁇ and massive microstructures are formed to cause a decrease in stretch flange formability. Thus, the average cooling rate (CR4) is limited to 0.01°C/s or more. When, on the other hand, the average cooling rate (CR4) exceeds 5°C/s, the partitioning of carbon to retained ⁇ is suppressed and a sufficient amount of carbon-enriched regions cannot be obtained. Furthermore, fresh martensite is formed to cause a decrease in ⁇ . Thus, the average cooling rate (CR4) is limited to 5°C/s or less.
  • the average cooling rate (CR4) is "(cooling start temperature (T3)) - (finish cooling temperature (T4))/(cooling time (seconds) from cooling start temperature (T3) to finish cooling temperature (T4))".
  • the cooling start temperature (T3) and the finish cooling temperature (T4) are not particularly limited as long as they are in the range of 340°C or above and 590°C or below.
  • the cooling start temperature (T3) is preferably in the range of 360 to 580°C.
  • the finish cooling temperature (T4) is preferably in the range of 350 to 450°C.
  • the holding (residence) in the temperature range of 340 to 590°C may also include a hot-dip galvanizing treatment. That is, the steel sheet may be subjected to a hot-dip galvanizing treatment or a hot-dip galvannealing treatment in the step where the steel sheet is caused to reside while being cooled at an average cooling rate (CR4) of 0.01 to 5°C/s.
  • a hot-dip galvanizing treatment is performed, the steel sheet is hot-dip galvanized by being immersed into a galvanizing bath at 440°C or above and 500°C or below, and the coating weight is preferably adjusted by, for example, gas wiping.
  • the galvanizing bath used in the hot-dip galvanization preferably has an Al content of 0.10% or more and 0.22% or less.
  • a hot-dip galvannealing treatment may be performed by an alloying treatment of the zinc coating after the hot-dip galvanizing treatment.
  • the alloying treatment of the zinc coating is preferably performed in a temperature range of 470°C or above and 590°C or below. Although this step is a cooling step (residence and slow cooling), the hot-dip galvanizing treatment and the alloying treatment of the zinc coating may be performed during the step as long as the temperature range, the residence time, and the average cooling rate CR4 described above are satisfied.
  • the hot-dip galvanizing treatment and the alloying treatment of the zinc coating may involve a temperature rise.
  • Cooling to a temperature of 50°C or below at an average cooling rate (CR5) of 0.1°C/s or more
  • the steel sheet is then cooled to a temperature of 50°C or below at an average cooling rate (CR5) of 0.1°C/s or more.
  • an average cooling rate (CR5) of 0.1°C/s or more.
  • the steel sheet may be subjected to skin pass rolling.
  • the skin pass rolling reduction ratio is preferably 0.1 to 0.5%.
  • the sheet shape may be flattened with a leveler.
  • the average cooling rate (CR5) to a temperature of 50°C or below is preferably 5°C/s or more, and more preferably 100°C/s or less.
  • the average cooling rate (CR5) is "(340°C (cooling start temperature) - finish cooling temperature of 50°C or below)/(cooling time (seconds) from cooling start temperature to finish cooling temperature)".
  • the above heat treatment or the skin pass rolling may be followed by a low-temperature heat treatment at 100 to 300°C for 30 seconds to 10 days.
  • This treatment tempers martensite that has been formed during the final cooling or the skin pass rolling and detaches from the steel sheet hydrogen that has penetrated into the steel sheet during annealing.
  • hydrogen can be reduced to less than 0.1 ppm.
  • electroplating may be performed. That is, the steel sheet may be electrogalvanized after the step where the steel sheet is cooled at an average cooling rate (CR5) of 0.1°C/s or more.
  • the electrogalvanization is preferably followed by the above low-temperature heat treatment in order to reduce the amount of hydrogen in the steel.
  • the steel sheet of the present invention preferably has a thickness of 0.5 mm or more.
  • the thickness is preferably 2.0 mm or less.
  • the member of the present invention is obtained by subjecting the steel sheet of the present invention to at least one working of forming and joining.
  • the method for manufacturing a member of the present invention includes a step of subjecting the steel sheet of the present invention to at least one working of forming and joining to produce a member.
  • the steel sheet of the present invention has a tensile strength of 980 MPa or more and is excellent in ductility, stretch flange formability, and laser weldability.
  • the member that is obtained using the steel sheet of the present invention also has high strength and has excellent ductility, excellent stretch flange formability, and excellent laser weldability compared to the conventional high-strength members.
  • weight can be reduced by using the member of the present invention.
  • the member of the present invention may be suitably used in an automobile body frame part.
  • the member of the present invention also includes a welded joint.
  • the forming may be performed using any common working process, such as press working, without limitation.
  • the joining may be performed using common welding, such as spot welding or arc welding, or, for example, riveting or crimping without limitation.
  • Steel sheets of the present invention and steel sheets of COMPARATIVE EXAMPLES were manufactured by treating 1.4 mm thick cold rolled steel sheets having a chemical composition described in Table 1 under annealing conditions described in Table 2.
  • the cold rolled steel sheets had been obtained by subjecting steel slabs having the chemical composition described in Table 1 to hot rolling (slab heating temperature: 1200°C, soaking time: 60 minutes, finish rolling temperature: 900°C, coiling temperature: 500°C) and cold rolling (rolling reduction ratio (cumulative rolling reduction ratio): 50%).
  • the martensite start temperature Ms was obtained using a Formaster tester by holding a cylindrical test piece (3 mm in diameter ⁇ 10 mm in height) at a predetermined annealing temperature and quenching the test piece with helium gas while measuring the volume change.
  • Some of the steel sheets (cold rolled steel sheets: CR) were obtained as hot-dip galvanized steel sheets (GI) by being hot-dip galvanized in a step where the steel sheet was caused to reside in a range of temperatures of 340°C or above and 590°C or below for 20 seconds or more and 3000 seconds or less while being cooled at an average cooling rate of 0.01 to 5°C/s.
  • the steel sheets were hot-dip galvanized by being immersed into a galvanizing bath at a temperature of 440°C or above and 500°C or below, and the coating weight was adjusted by, for example, gas wiping.
  • the galvanizing bath used in the hot-dip galvanization had an Al content of 0.10% or more and 0.22% or less.
  • hot-dip galvanized steel sheets were alloyed and obtained as hot-dip galvannealed steel sheets (GA) by being subjected to an alloying treatment after the hot-dip galvanizing treatment.
  • the alloying treatment was performed in a range of temperatures of 460°C or above and 590°C or below.
  • steel sheets cold rolled steel sheets: CR
  • electrogalvanized steel sheets EG
  • the steel microstructure was measured in the following manner. The measurement results are described in Table 3.
  • the steel sheet was cut to expose a through-thickness cross section that was parallel to the rolling direction.
  • the cross section was mirror-polished and was etched with 3 vol% Nital. Portions at 1/4 thickness were observed with SEM in 10 fields of view at a magnification of 5000 times.
  • the polygonal ferrite discussed here is relatively equiaxed ferrite containing almost no internal carbides. This region looks blackest in SEM.
  • Bainitic ferrite is a ferrite microstructure that contains carbide or retained ⁇ which looks white in SEM.
  • the area fractions were calculated while assuming that the region was polygonal ferrite when the aspect ratio was ⁇ 2.0 and the region was bainitic ferrite when the aspect ratio was > 2.0.
  • the aspect ratio was obtained from a/b in which a was the major axis length where the particle length was longest, and b was the minor axis length that cut across the particle over the largest length perpendicular to the major axis length.
  • the particles were divided at a position that divided the particles approximately evenly and the sizes of the respective particles were measured.
  • the number of internal carbides in bainitic ferrite per 10 ⁇ m 2 was determined by measuring the area fraction of bainitic ferrite and counting the number of internal carbides in a ⁇ 5000 SEM image, dividing the carbide count by the area fraction of bainitic ferrite, and converting the quotient into the value per 10 ⁇ m 2 .
  • Tempered martensite is a region that contains a lathlike submicrostructure and carbide precipitates according to SEM.
  • Fresh martensite is a massive region that looks white and does not contain any visible submicrostructures according to SEM.
  • the microstructure including one, or two or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite corresponds to the remaining microstructure after the polygonal ferrite, and the total area fraction of this microstructure is the area fraction of the regions other than the polygonal ferrite.
  • the area fraction of carbides was very small and was thus included in the above area fraction of the remaining microstructure.
  • the volume fraction of retained austenite (retained ⁇ ) was determined by chemically polishing the steel sheet surface to a location at 1/4 thickness and analyzing the sheet surface by X-ray diffractometry. Co-K ⁇ radiation source was used as the incident X-ray, and the volume fraction of retained austenite was calculated from the intensity ratio of (200), (211), and (220) planes of ferrite and (200), (220), and (311) planes of austenite. Because retained ⁇ is randomly distributed, the volume fraction of retained ⁇ obtained by X-ray diffractometry is equal to the area fraction of retained ⁇ in the steel microstructure.
  • the number density of retained austenite present adjacent to bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 was determined as follows. The sample used for bainitic ferrite observation was mirror-polished. An electron backscattering diffraction pattern (EBSD) of the same field of view as obtained in SEM was subjected to mapping measurement using EBSD analysis program OIM Data Collection ver. 7. The obtained data was analyzed using TSL OIM Analysis ver. 7 (manufactured by EDAX/TSL) to give phase map data. The number density of fcc structures present adjacent to bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 was measured.
  • EBSD electron backscattering diffraction pattern
  • adjacent means that the fcc structures were in contact with bcc structures corresponding to bainitic ferrite having 20 or less internal carbides per 10 ⁇ m 2 in the phase map.
  • the fcc structure may be included in the bcc structure.
  • the area S C ⁇ 0.5 of a region having a C concentration of 0.50% (mass%) or more and the area S C ⁇ 0.3 of a region having a C concentration of 0.30% (mass%) or more were measured by mapping analysis of the C concentration distribution with respect to positions at 1/4 thickness of a through-thickness cross section perpendicular to the steel sheet surface and parallel to the rolling direction, using field emission electron probe microanalyzer (FE-EPMA) JXA-8500F manufactured by JEOL Ltd., at an acceleration voltage of 6 kV and an illumination current of 7 ⁇ 10 -8 A with the minimum beam diameter.
  • FE-EPMA field emission electron probe microanalyzer
  • JXA-8500F manufactured by JEOL Ltd.
  • JIS No. 5 test pieces for tensile test and test pieces for hole expansion test were sampled from the steel sheets obtained.
  • a tensile test was performed (in accordance with JIS Z2241).
  • the tensile strength TS and the total elongation T-El are described in Table 3.
  • the steel sheets were evaluated as being excellent in strength when the tensile strength was 980 MPa or more.
  • the ductility was evaluated as excellent when the total elongation T-El was 16.0% or more for the steel sheets having a TS of less than 1180 MPa, when the total elongation T-El was 14.0% or more for the steel sheets having a TS of 1180 MPa or more and less than 1320 MPa, and when the total elongation T-El was 13.0% or more for the steel sheets having a TS of 1320 MPa or more.
  • test pieces for hole expansion test sampled from the steel sheets after the heat treatment were subjected to a hole expansion test conforming to the provisions of The Japan Iron and Steel Federation Standard JFST 1001 to evaluate stretch flange formability.
  • a 100 mm ⁇ 100 mm square sample was punched with a punching tool having a punch diameter of 10 mm and a die diameter of 10.3 mm (13% clearance), and a conical punch having an apex angle of 60 degrees was inserted into the hole in such a manner that the burr produced at the time of punching would be directed to the outside.
  • the hole was expanded until the sheet was cracked through the thickness.
  • the hole expansion ratio ⁇ (%) ⁇ (d - d 0 )/d 0 ⁇ ⁇ 100 was calculated.
  • d 0 initial hole diameter (mm)
  • d hole diameter (mm) at the occurrence of cracking.
  • Table 3 The steel was evaluated as having excellent flangeability when ⁇ was 30% or more.
  • the fracture mode was determined to be the weld fracture. Furthermore, a test piece was sampled from the welded member in such a manner that the weld line would be perpendicular to the tensile axis and would be located at the longitudinal center of the test piece, and the weld was notched as illustrated in Fig. 1(b) . The notched test piece was subjected to a tensile test (a notched tensile test).
  • the test pieces were deformed exclusively at the HAZ of the weld and limited regions around the HAZ and were forcibly fractured at the HAZ portion to evaluate the strength of the HAZ portion itself.
  • the laser weldability (the HAZ softening resistance) was evaluated as being excellent when the test piece was fractured in the base material fracture mode in the fracture mode determination test and satisfied HAZ strength ⁇ base material TS + 50 MPa in the notched tensile test.
  • the steel sheets of the present invention have superior ductility, excellent stretch flange formability, and excellent laser weldability and can be suitably applied to press forming and be suitably used in the press forming process in the manufacturing of, for example, automobiles and home appliances.
  • the steel sheets of INVENTIVE EXAMPLES have high strength, excellent ductility, excellent stretch flange formability, and excellent laser weldability. This has shown that members obtained by forming of the steel sheets of INVENTIVE EXAMPLES, members obtained by joining of the steel sheets of INVENTIVE EXAMPLES, and members obtained by forming and joining of the steel sheets of INVENTIVE EXAMPLES will have high strength, excellent ductility, excellent stretch flange formability, and excellent laser weldability similar to the steel sheets of INVENTIVE EXAMPLES.

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EP22864236.9A 2021-08-31 2022-08-15 Stahlblech, element und verfahren zur herstellung des stahlblechs und des besagten elements Pending EP4361304A4 (de)

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JP5463685B2 (ja) 2009-02-25 2014-04-09 Ｊｆｅスチール株式会社 加工性および耐衝撃性に優れた高強度冷延鋼板およびその製造方法
JP5780086B2 (ja) 2011-09-27 2015-09-16 Ｊｆｅスチール株式会社 高強度鋼板およびその製造方法
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JP6472692B2 (ja) * 2015-03-23 2019-02-20 株式会社神戸製鋼所 成形性に優れた高強度鋼板
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WO2017208759A1 (ja) * 2016-05-30 2017-12-07 株式会社神戸製鋼所 高強度鋼板およびその製造方法
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Ipc: C21D 8/02 20060101AFI20240904BHEP